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By J. B. Cheatham, R. B. Paslay

Rock stresses and steady-state flow rates induced by the pressure gradient associated with the flow of formation fluid into a borehole have been analytically determined for a permeable, elastic material saturated with an incompressible fluid. In this analysis, the material properties and loading are considered to be symmetric about the axis of the borehole and independent of axial position. For Case I the material is assumed to have uniform permeability in the radial direction, whereas for Case 11 the permeability is assumed to have been reduced in a localized region adjacent to the hole by either normal well completion and production operations or deliberate plugging during air drilling. Results of a numerical example indicate that, in the absence of plugging, the rock shear strength must be approximately two-thirds the formation fluid pressure in order to prevent rock failure. The required rock strength is high for small radial zones of plugging and decreases as the region of reduced permeability becomes larger; however, a depth of plugging can be reached beyond which there is no real gain in strength, although the flow rate can be further reduced. INTRODUCTION During normal production of oil from a well, it is often desirable to increase the production rate of the formation fluid by increasing the pressure gradient through the formation adjacent to the borehole. Depending upon the magnitude of this pressure gradient and strength of the rock material, this production-rate increase can cause sloughing of the hole wall. In many cases, the production - rate increase can result in excessive sand production, increased wear of production equipment, lost production time and expensive workover jobs. In addition, the phenomenon of increased rock bit penetration rate with the use of a gaseous instead of a liquid drilling fluid has been observed in oilfield drilling operations and experimentally demonstrated by various investigators for several years. The improvement obtained by employing this technique can be quite significant and offers apromising method for reducing drilling costs. However, air drilling is currently limited to geographical locations where high-capacity water-bearing formations are not encountered. This limitation has prevented widespread adoption of air-drilling techniques, because the water influx into the borehole interferes with efficient removal of the drilling cuttings and usually results in a condition such that the bit becomes "balled-up" or stuck in the hole. In an attempt to remove the water - intrusion limitation from air drilling, various chemical and mechanical water shut-off methods have been proposed. Goodwin and Teplitzl suggested one such proposal whereby the permeability of the water-bearing rock structure was reduced in the vicinity of the borehole. Although the development of a shut-off method based upon this approach would certainly be welcomed by the oil industry, it is conceivable that, under certain conditions of the pressure gradient, strength of the rock material and depth of the modified permeability zone, a stress field can be created that will result in an unstable hole. As part of their study, an analytical solution is given for stresses in an idealized model of a hole and the surrounding rock material. The purpose of the present study is to extend the analysis of Goodwin and Teplitz to gain more insight into the details and consequences of excessive production rates and formation water shut-off. In particular, simplified models of these problems have been analytically examined, which makes possible the evaluation of the type of stress fields that can be anticipated as a result of these production and drilling practices. Both problems solved concern the determination of the steady-state volume flow rate of the formation fluid and the resulting steady-state stress and displacement distribution in a hollow, cylindrical configuration. The cylinder of Case I, corresponding to the production-rate problem, consists of a material with a constant permeability from the inside surface to the outside surface; the cylinder of Case 11, corresponding to the water shut-off problem, consists of a material with a constant permeability from the inside surface to an intermediate concentric cylindrical surface and a second constant permeability from the intermediate surface to the outside surface-

By A. E. Dwight

Lattice parameters were determined for eighteen equiatornic alloys of the CsCl-type structure, ten of which were previously un-reported. It was found that fomation of the CsCl-type structure in binary alloys of the transition elements is largely dependent on position of the elements in the periodic table. The relative size of the two elements was not found to be a controlling factor. A recent paper by Beck, Darby, and Arora1 corre-lates the occurrence of CsC1-type ordered structures with the position of the constituent elements in the periodic table for the first long period. It was also suggested that a definite increase in relative bond strength between unlike atoms occurred when, in binary alloys of iron-group elements, the other component is changed from a chromium-group element, to a vanadium-group element, to titanium. A later paper by Philip and Beck2 noted that the lattice contraction increased in the order CrFe, VFe, and TiFe. It was also noted by Philip and Beck2 that the lattice contractions of CsC1-type alloys decreased in the order: TiFe, TiCo, and TiNi, which is an apparent reversal of the contractions expected from the position in the periodic table. It was suggested that the increasing lattice contraction is an indication of increased stability, i.e., greater A-Bbond strength. The present investigation was carried out to determine whether the relation of the position in the periodic table to the formation of the CsC1-type structure was also correct for alloys involving the second and third long-period elements. A systematic search was made for CsC1-type structures among equiatomic alloys and for those found, the lattice contraction was determined. EXPERIMENTAL TECHNIQUE The elements Y, Gd, Ti, Zr, Hf, V, and Cb are designated the A group and the elements Mn, Re, Fe, Ru, Os, Co, Rh, Ir, Ni, Pd, Pt, Cu, Ag, and Au are designated the B group. Equiatomic alloys were prepared for 57 AB combinations. The alloys were arc melted in a multicrucible furnace3 in buttons ranging from 5 to 20 g. Chemical analyses were not made, as the charge weights agreed closely with those of the buttons after melting. The alloy buttons were homogenized at 800° to 900°C. Metal log raphic and X-ray specimens were prepared and heat treated at temperatures from 600° to 1200°C. Specimens for X-ray diffraction were usually ground to a powder in an agate mortar; however, needle-shaped solid specimens were used when the alloy was sufficiently ductile to permit their preparation. Diffraction patterns were taken with a Straumanistype Debye-Scherrer camera using filtered Cu or Co radiation. The lattice parameters were obtained in A by plotting the calculated a0 values against the cos29/sin 0 + cos2?/? function and extrapolating linearly to ?= 900. Metallographic control specimens were polished on cloth wheels with diamond paste and etched with various phosphoric and nitric acid reagents. RESULTS The eighteen equiatomic alloys listed in Table I gave evidence of a cubic structure with two atoms in the unit cell, although two of these cubic structures exist only at elevated temperatures and transform to a tetragonal structure on quenching. Nine of these eighteen alloys gave diffraction patterns with super-lattice lines showing that the structure is of the CsC1-type. The lack of superlattice lines in patterns of the other nine alloys may be attributed to the small difference in atomic scattering power of the components. Metallographic study indicates the occurrence of nine narrow single-phase fields at the AB composition. Any or all of these nine may also have a CsC1-type structure. The VFe alloy was found to have a CsCl-type structure by Philip and Beck2 through the use of CrKa radiation (for which the scattering factor of V is anomalously low), whereas the Cu radiation used

Jan 1, 1960

By F. D. Rosi, C. A. Dube, B. H. Alexander

The slip and twinning planes have been determined in deformed crystalsof titanium by an X-ray method of analysis. The slip planes are of the type {1010} and {1011}, while the twinning planes are of the type {1072}, {1121}, and (1122). In the case of the predominant (1010) slip, a type I digonal axis of indices <1120> was the effective slip direction. Microscopic observations on the appearance of the three twin types disclosed a distinct difference in their shapes, and an attempt was made to correlate this difference with the amount of twinning shear. The slip mechanism was discussed on the basis of the difference in c/a ratios for the common hexagonal close-packed metals. IT is a general rule that slip in metals occurs most easily on the planes of greatest atomic density and of the largest interplanar spacing, and that the direction of slip is a close-packed direction. Thus, of the hexagonal close-packed metals which have been investigated by the German school in the decade 1925 to 1935 (cadmium, zinc, and magnesium), it has been found that slip occurs on the basal plane (0001) and in a <1120> direction.&apos; In the case of magnesium above 225oC, slip occurs on a first-order pyramidal plane (1011) as well as on the basal plane.&apos; The hexagonal metals, moreover, deform in part by twinning, and, again with the exception of magnesium at elevated temperatures, the twinning plane has always been reported to be a second-order pyramidal plane {1012).&apos; In view of this consistency in the plastic behavior of these metals, it has been supposed that the same slip and twinning elements would be operative in the other important hexagonal close-packed metals (titanium, zirconium, and beryllium), although there are some reasons why this may not be so. Referring to the tabulation in Table I of the axial ratios of the important metals in this structural class, it may be seen that they can be categorized into three distinct groups depending on the value of their c/a ratios. In the first group are cadmium and zinc with large positive deviations from the ideal ratio for closest packing. This expansion in the direction of the c-axis would be such as to accentuate the basal plane as the slip plane, because it increases the interplanar spacing. The second group includes cobalt and magnesium, which closely approximate the ideal ratio. In the third group are zirconium, titanium, and beryllium whose axial ratios are decidedly less than the ideal. The lattice of the elements in this group is compressed along the c-axis which, in effect, tends to make the basal plane less favorable for slip inasmuch as it reduces the interplanar spacing and the atomic density of this plane. Thus, for metals in the third group, it may be said that there is a certain indefiniteness about the glide plane, in that planes other than the basal plane may be expected to operate. The fact that magnesium begins to exhibit planes other than the normal slip and twinning planes at elevated temperatures may indicate that this behavior will become more pronounced in metals of lower axial ratios. For this reason it might be expected that the slip and twinning behavior of zirconium, titanium, and beryllium would be different from that found in cadmium and zinc. One indication that there is a difference is reflected in the recent work which has been done on the deformation textures of these metals."&apos; As will be shown later, there is a significant difference between the rolling textures of cadmium, zinc, and magnesium on the one hand, and zirconium, titanium, and beryllium on the other, and this difference suggests that the mechanism of plastic flow is not the same in these two groups of metals. From these considerations, therefore, it may be seen that there are good reasons why it cannot be concluded that titanium will have the same crystal-lographic elements of slip and twinning as those exhibited by zinc and cadmium. For the complete determination of these elements it would be most advantageous to have single crystals. However, much valuable information can be obtained by using coarse grained polycrystalline specimens, provided the grains are sufficiently large to permit X-ray

Jan 1, 1954

By C. D. King

DURING the 30-year period spanned by these annual Howe Memorial presentations, many lecturers could proudly claim a kinship either as a student or an associate of the man whose memory we honor. Although it has been my good fortune to have attended many of these annual lectures, it was not my privilege to have known Henry Marion Howe personally. However, his great repute as teacher and scientist was known to all undergraduates of my day and the later years have enhanced my appreciation of his wisdom and foresight. Those who knew him well have said he derived particular pleasure from speculations on the future world of metallurgy. For this reason, I feel that perhaps he would not be unsympathetic to a lecture in his honor which departs from the highly instructive scientific presentations made in the past by so many able Howe Memorial lecturers, and which is concerned more with the practical phases of various steelmak-ing processes and some speculations on their future form and relative importance. The word "revolutionary" is frequently applied to each seemingly important improvement in the production of steel ingots, but in retrospect these changes, impressive as they appear at the time, are merely steps of progress. In the hundred years from the inception of tonnage steelmaking, only three processes can be truly classified as revolutionary. They are the pneumatic process, known in this country as the bessemer process; the reverberatory method called the open hearth process; and, the electric furnace process. There have been many variations and combinations of the three fundamental methods, but they remain truly the only revolutionary methods in steelmaking since its early history. Everything else has been evolutionary, in effect. doing the same things that we have done in the past but doing them better, correcting our errors through experience, and slowly but inevitably reaching a higher state of accomplishment. It has often been said that coming events cast their shadows before, and the production of steel ingots is no exception. As a result of the unrelenting demands of World War I1 and the years that followed, truly impressive progress has been made in steel ingot production. The incessant pressure for immediate results during this period required the employment of initiative and daring, as in few past decades, and many developments were brought to fruition. Of equal importance is the possible effect on future steelmaking methods of the many ideas initiated but still in formative stages. Fig. 1 portrays ingot production in the United States by the three fundamental processes over a period of 75 years and is interesting because it poses some questions as to future trends. The early ascendancy of the bessemer, its replacement in importance by the open hearth process, the amazing growth of the latter, and the recent challenge of the electric furnace are evident from the chart. Management is fully aware of these changes, but is even more interested in the future trends. Our concepts of the relative importance of the more recent developments and their possible effect on future processes may perhaps be best exemplified by a specific, hypothetical problem. Let us assume management is contemplating a new ingot producing plant with an output of 100,000 net tons per month, located in an area where some purchased scrap may be obtained but where by far the largest component will be own-produced blast furnace iron. Management requires a process or combination of processes which will yield highly uniform quality characteristics in the ingot form, and represent the soundest selection in investment and operating cost. Under these conditions, the obvious selection for the past four decades has been the open hearth process but, in view of more recent developments, management may believe that it is no longer permissible to disregard other possibilities with impunity. Accordingly, to be assured of the best possible selection, they request that you review not only the possibilities of utilizing the conventional open hearth, duplex, bessemer, and electric furnace methods, but also the more recent developments, such as the turbo-hearth, the Linz-Donawitz method, the Perrin modifications, and other possibilities. With this background, one might then appraise the relative importance of these methods to meet a specific need, and concurrently speculate on the forms that future ingot processes will assume and the relative importance of these processes.

Jan 1, 1955

By R. Stickler, A. Vinckier

An optical and electron microscope study was made on a commercial 302 stainless steel heat treated for massive carbide precipitation. Convincirg evidence was obtained that the grain boundary precipitate does not grow away from the grain boundary into the matrix, but grows as a network of large thin dendrites in between the grains. IN austenitic stainless steels the carbon in supersaturated solution precipitates when the steel is heated in the temperature range 450" to 1000°C. During such heat treatments the chromium-rich carbide occurs primarily at the grain boundaries and has adverse effects on the corrosion resistance and low-temperature ductility of the material. It has been shown that the morphology of the precipitated particles differs widely depending on the time and temperature of the heat treatment, the nature of the boundary, the misfit between the grains, and so forth.1"5 It was assumed by Mahla et al.,&apos; Streicher,&apos; and Mc Nutt&apos; that the carbide particles which nucleate in the grain boundary grow in various geometric or dendritic forms away from the boundary into the matrix. However, Kinzel,&apos; Hatwell et a, and Plateau et aL4 found that the particles nucleate in the grain boundary and grow as a network of thin dendrites in the grain boundary interface. We have made a comprehensive studyg of the precipitation occurring in a commercial 302 austenitic stainless steel heat treated for times ranging from 0.15 to 1500 hr at various temperatures from 480" to 1065"~. This study was made to elucidate the influence of heat treatments on both the mechanical and the corrosion properties of this steel. Part of this investigation which is pertinent to the morphology of the grain boundary carbides referred to above is separately reported here. We found that the particles which nucleated in the boundaries grew in the interface of these boundaries, thus supporting the viewpoints of Kinzel,&apos; Hatwell et al.,3 and Plateau et al.4 Furthermore, when carbide particles nucleated in the matrix very close to the boundaries, they possessed a morphology distinctively different from the particles growing in the grain boundaries. MATERIAL A study of metallographic samples and of fracture surfaces of a 302 stainless steel is reported in this paper. The composition and heat treatments are listed in Table I. Polished and etched samples were examined und813r the optical microscope, and carbon extraction replicas of etched micrographic samples and various fracture surfaces3&apos;8&apos;9 were examined under the electron microscope. RESULTS AND DISCUSSION A representative optical micrograph of the 302 steel, condition A, is shown in Fig. l(a). Particles can be seen on the grain boundaries but no traces can be found of particles which grow from the grain boundaries into the matrix. A carbon extraction replica of this sample shows numerous large thin dendrites, Fig. l(b). The lateral dimension of such carbide particles vary up to 100 p, and their thickness estimated from the amoun! of electron penetration i:; between 500 and 1000A. Such large dendrites are extracted from almost all grain boundaries and, if they had grown into the grains, they would definitely be revealed by etching traces in the matrix, Fig. l(a). The reason why these large dendrites on the extraction replica appear to be grown into the grain is due to the mechanism of the extraction replica process. Thus, these dendrites are supported by the carbon replica film only at the edges originally along the grain boundary trace A-A, Fig. l(b), in the polished metal section and have subsequently fallen over onto the carbon replica film during handling. Although such dendrites generally fall only to one particular side of the grain boundary trace, one finds occasionally that parts of a dendrite have fallen to both sides, as shown in Fig. l(b). Further evidence for the growth of carbide in the grain boundary can be obtained from the appearance of fracture surfaces. A small impact specimen of the 302 steel, condition A, was broken at liquid nitrogen temperature. The fracture path follows the

Jan 1, 1962

By J. R. Hauber, N. R. Borch

A study uras made on the creep behavior of cast and extruded SR grade beryllium. It is shown that, for stresses below about 1000 psi in the temperature range 760" to 85o° c, the creep behatior is nearly exactly described by the Nabarro-Herring mechanism. The activation energy is apparently that for self-diffusion, 42 2 kcal per mole, below 1000 psi, but it rises to 88.6&apos; kcal per mole at higher stresses. The slress exponent is 4.5 at the higher stresses. Tertiary creep in this material is clearly related to void formation apparently caused by grain boundary sliding. BERYLLIUM is becoming an increasingly important technological material. A search of the literature has revealed a lack of fundamental studies on the creep of beryllium. This investigation was therefore undertaken to gain insight into the creep behavior of this material. EXPERIMENTAL PROCEDURE Tensile creep samples were prepared from cast and extruded SR grade beryllium. The vendor&apos;s chemical analysis of the beryllium tested is shown in Table I. The test specimens were machined as shown in Fig. 1 from i -in.-diam rods. The machined creep specimens were annealed to remove surface damage and to establish a variation in grain size. The heat treatments and their resulting grain diameters are shown in Table 11. The average grain diameters were determined by the intercept method. True average grain diameters were estimated by multiplying the measured average grain diameters by 1.5.&apos; The creep tests were performed at temperatures from 700" to 850°C in a vacuum of 105 Torr or better in a tantalum-element resistance furnace. Loads were applied by an Instron testing machine such that the stress was held constant to within 20 psi. The test temperature was measured by thermocouples on the sample. With one exception the deviation from the desired test temperature was 2"C, and in that case the deviation was 4"C. Elongations were measured by an extensometer within the vacuum chamber which was fastened to the creep specimens at the grooves in the grip ends, as shown in Fig. 1. The precision of measurement of elongation was 1 X loe4 in. The creep tests were begun only after thermal equilibrium had been established in the sample and the extensometer to insure that there was no strain rate component due to transient thermal effects. Two kinds of creep experiments were performed. First, at constant grain size, the effects of varying stress and temperature were studied. In the second set of experiments, the temperature was held constant and the grain size was varied. The activation energy for creep was determined by varying the temperature at constant stress, as described by orn, during the first set of experiments. The variable grain size experiments revealed that creep at low stress apparently occurred by the abarro-errin&apos; mechanism. In all of the creep tests, the stress was held constant only until sufficient strain had occurred to accurately determine the creep rate. The stress was then raised to a new value and the new creep rate determined. This process was reported until tertiary creep occurred. In this manner several data points were obtained from each specimen. The assumption was made that the structure was not changing appreciably between the incremental stress levels. This was verified by reducing stresses to previously tested levels and comparing the creep rates with the creep rates obtained at the lower strains. The re-producibility was within experimental error over strains on the order of 1 pct. EXPERIMENTAL RESULTS A typical series of creep tests is shown in Fig. 2. Primary creep was not observed in any test. Transient creep, which we define as creep rates that are nonlinear with time, was not observed for any change in stress except where tertiary creep occurred. The absence of both primary and transient creep has been previously reported for hot-pressed powder beryllium.= The data for the creep behavior at constant grain size are summarized in Fig. 3. The data show two distinctly different stress dependencies. At stresses below about 1000 psi, in the equation:

Jan 1, 1969

By E. S. Machlin, Morris Cohen

The martensite reaction in single crystals and polycrystals of 70 pct Fe-30 pct Ni alloys is shown to be autocatalytic in nature, producing bursts of transformation during cooling. The temperature of the first burst of transformation, called Mb, occurs below M, in these alloys. Experiments were devised to test the athermal embryo and strain embryo theories of martensite nucleation. The results indicate that internal strains, either within the virgin austenite or around existing martensitic plates, control the nucleation process in these alloys. Furthermore, the growth of martensitic plates is not limited by the attainment of an elastic balance with the austenitic matrix, but by the occurrence of plastic deformation at the martensite boundaries which interferes with the propagation mechanism. IN an investigation of the martensitic habit in single crystals of a 69 pct Fe-31 pct Ni alloy,&apos; it was observed that about 25 pct of the austenite transformed during subatmospheric cooling within the time-interval of an audible click. This event proved quite spectacular: The shock wave sent out from the specimen freely suspended on a thread in the refrigerating liquid was occasionally sufficiently intense to shatter the Dewar container and to separate the toluene column in the immersed thermometer. The Present investigation was undertaken to determine- the kinetics and mechanism of this "burst" type of martensitic reaction. The analyses of the alloys studied are given in Table I. The composition of the single crystal specimens is designated by alloy A, while the polycrystal-line specimens were made of alloys B and C as noted in the text. The single crystals were prepared in a vacuum furnace, using a modified Bridgman technique. Most of these crystals were homogenized by holding for 24 hr at about 1300°C just after solidification. However, it may be emphasized here that the degree of homogenization was not a controlling factor in the subsequent experiments, inasmuch as specimens having different degrees of homogenization yielded the same results. All of the single crystals were fully austenitic as slowly cooled to room temperature. An illustration of the burst phenomenon is given in Fig. 1, which shows oscillograms of electrical resistivity and temperature vs. time during the continuous quenching of 1/16 in. wire specimens (alloy B) in a dry ice and acetone bath at —77°C. There are at least two observable bursts in this case, as indicated by the sharp decreases of resistance accompanying the sudden formation of substantial quantities of martensite. The thermal arrest during the quench probably corresponds to the larger burst. Usually the bursts are followed by more or less progressive transformation during continuous cooling. It will also be noted that the resistance continues to decrease after the specimen has reached the bath temperature. This isothermal change denotes the formation of martensite at constant temperature, and will be the subject of another paper. Examination of fiducial scratches on the surface of a transformed single crystal has shown2 that the scratches in adjoining nonparallel martensitic plates are usually bent in opposite directions, as though one plate forms in such a way as to relieve the matrix stresses set up by the adjacent plate. This, together with some of the results described in ref. 1, Table I. Compositions of Alloys Studied, in Percent Alloy Ni C N Mn Si P S Cr A 31±0.3 0.048 0.027 0.003 0.56 0.007 0.002 B 29.5±0.2 0.036 0.02 0.19 0.09 0.008 0.006 C 19.99 0.52 0.37 0.47 0.010 0.015 0.04 led to the tentative concept that a cooperative action exists which provides the impetus for much of the transformation that appears during the burst. The following series of experiments were performed in order to test this idea. Cooperative Nature of the Burst Two adjacent disks, Va in. thick x % in. diam, were cut from a single austenite crystal of alloy A using a jeweler&apos;s saw. One of the disks was then cut into 15 parts. Then 12 of the latter pieces and the second disk were austenitized (stress relieved) at 600°C for 30 min and water quenched to room temperature. The temperatures at which the first burst of transformation appeared were determined for

Jan 1, 1952

By T. W. Watson, D. R. Flynn

Measurements have been made of the thermal conductivity and electrical resistivity of two specimens of 0.005-cm (2-mil) Be-Cufoil over the temperature range -140° to +200°C. The thermal cmductivity of Be-Cu Alloy 125 was found to be significantly higher than that of Alloy 25, which had a higher impurity content. The thermal conductivity and electrical resistivity values obtained were in concordance with the Smith and Palmer relation for copper alloys which states that ? = 0.0239 (T/p) + 0.075 where A is thermal conductivity (W cm-&apos; deg-&apos;), p is electrical resistivity (microhm-cm), and T is absolute temperature (OK) , indicating that the Smith and Palmer relation can be used to predict the thermal conductivity of this type of Be-Cu alloy over the temperature range -140" to + 200°C to within an accuracy suitable for most engineering applications. THE National Aeronautics and Space Administration, Goddard Space Flight Center, Greenbelt, Md., requested the National Bureau of Standards to measure the thermal conductivity of 0.005-cm Be-Cu foil which was to be used to fabricate antennae for use in Radio Astronomy Explorer space craft. It was felt that measurements had to be made on the actual antenna material, rather than on a thick bar of similar material, since the cold-working and age-hardening properties of the alloy would make it very difficult, if not impossible, to place a thick bar in the same metallurgical state as the 0.005-cm material. As discussed in Section 111 of this paper, the thermal conductivity of many copper alloys can be computed rather accurately from electrical resistivity data using an empirical relation developed by Smith and palmer&apos; and further confirmed by owell.&apos; However, there has been very little work done on correlating thermal and electrical conductivities on Be-Cu alloys. We were reluctant to assume the validity of the Smith-Palmer relation on Be-Cu alloys because of the very large mass difference between beryllium and copper. As will be seen, however, our experimental results indicate that the Smith-Palmer relation is valid for the two alloys on which we made measurements. I) DESCRIPTION OF SAMPLES Samples of Be-Cu Alloy 25 and Be-Cu Alloy 125 were furnished to NBS in the form of rolled strip material, 0.005 cm thick by 5.08 cm wide. The Be-Cu Alloy 25 was purchased by NBS in the form of a continuous length of flat strip. The Be-Cu Alloy 125 was furnished to NBS by Goddard Space Flight Center in the form of two lengths of strip that had been formed into "tubing" of about 1.2 cm diam, with one side of the strip overlapping (but not connected to) the other side by about 90 deg. This "tubing" had been opened up to be flat and then rolled onto spools of about 3 cm diam. When unrolled, the strip would spring back into tubular form. Samples of both of the Be-Cu alloys were chemically analyzed by the NBS Analytical Chemistry Division. The beryllium content was determined quantitatively by the phosphate-gravimetric method to be 1.83 wt pct for Alloy 25 and 1.80 wt pct for Alloy 125. The two alloys were also examined by a general qualitative spectrochemical method for metallic elements. The Alloy 25 specimen contained between 0.1 and 1.0 wt pct Al, Co, Fe, and Si and between 0.01 and 0.1 pct Ni. The Alloy 125 specimen contained between 0.1 and 1.0 pct Co and between 0.01 and 0.1 pct Al, Fe, and Si. Other metallic elements present were in amounts less than 0.01 pct. The NBS Engineering Metallurgy Section measured the tensile strength, elongation, and Knoop hardness of both the Be-Cu Alloy 25 and the Be-Cu Alloy 125. These are given in Table I. The NBS Engineering Metallurgy Section also prepared photomicrographs of the two alloys; these are shown in Figs. 1 and 2.

Jan 1, 1969

By J. W. Snavely

THE part that acceleration plays in starting a belt conveyor and its effect on belt conveyor design are well understood in a general way. Its practical importance is easily overlooked, however, and under some conditions, it is absolutely necessary to give the problem of acceleration detailed study. Most handbooks on conveyor belting design adequately present basic data for the determination of acceleration values. This paper will only attempt to present practical thinking and a convenient method of treatment of acceleration in belt conveyor design. Mathematical Analysis In working out the various problems of conveyor belt acceleration, the starting point, as presented by the handbooks, is the familiar formula of "force of acceleration is equal to the mass times acceleration." By expressing these fundamental quantities in terms of belt conveyor design, it is possible to arrive at the unsuspected conclusion that the acceleration time for horizontal belt conveyors is independent of the load, and instead, dependent upon the belt speed, the type of drive arrangement and drive pulley, and the idler coefficient of friction. The mathematics leading to this conclusion are shown in Table I, which has been prepared to show, this derivation. While at first the conclusion just given may not seem to be reasonable, further reflection indicates that obviously the type of drive pulley and the type of drive do affect materially the tension in the conveyor belt, and thus, as clearly shown, the time of acceleration is dependent upon the factors mentioned. Inasmuch as all of the factors except time are predetermined by the belt conveyor design, it becomes relatively easy to establish the accelerating time and to reduce further this time determination to a simple graph from which the time in seconds can be read directly. Such a graph is given in Fig. 1. The table appearing on Fig. 1 should be explained further. For a given belt speed, the time of acceleration can be expressed as a percentage of the belt speed. The time of acceleration is also dependent on the drive arrangement, and changes in the drive arrangement consequently change the time of acceleration. It further follows that for a given belt speed, the time expressed as a percentage of that belt speed also changes with the type of drive. Obviously then, it becomes possible to graph the percentage of speed for each type of drive against the belt speed and accelerating time, after which, for a given belt speed and type of drive, the time can be read directly in seconds. Two constants were established for Fig. 1, the first one being the limiting of the maximum acceleration tension to 35 pct of the full load operating tension in the belt. The purpose of this is to limit the total tension imposed upon the belt during the acceleration period to 135 pct of the full load operating tension, which is the amount required to start or breakaway the fully loaded belt conveyor from rest. The other constant is the friction factor used for the idler equipment, which has been established as 0.022. For installations where it is necessary to establish the values of acceleration, invariably high grade idler equipment is used, and it has been established from. field experience that 0.022 for the idler friction factor is amply conservative. The use of this friction factor for idlers must be tempered with judgment, of course, for occasions will arise where more power than indicated is required to start, even with the very best of equipment, such as low temperature operations that tend to congeal the grease in the bearings and thus produce additional friction drag. An inspection of the table in Fig. I affords a convenient rule of thumb method for determining the acceleration time, which conveniently can be 5 pct of the belt speed in seconds. The 5 pct of belt speed figure is close to the average for most types of drives. In using Fig. 1 it must be emphasized that it applies accurately to horizontal belt conveyors only.

Jan 1, 1952

By Cloyd M. Smith

B. F. TiLLson*— The manifold difficulties of accurate anemometry in irregular sections of mine passageways, the irregular distributions of velocities in cross sections of the same, and the disturbing influence of the observer upon the air flow, all indicate an undesirable inaccuracy of results obtainable by any standardized method of traversing the section by an anemometer. It seems obvious that another, and simpler, method should be used to determine the volume of air flow in mine passages, namely: 1. At appropriate locations cement or calked framework rings should be installed permanently to equalize the irregularities of sectional contour and provide a place and means of attachment for a temporary cloth brattice which bears a rigid orifice. 2. The measurement of the velocity of air flow through the orifice may then be by anemometer or, preferably, by Pitot tube measurements of the differential pressures on both sides of the orifice in accordance with the standard practices available in engineering 1iterature. † The constants may be determined for various measuring positions in relation to the resulting "vena contracta." 3. The position of the person who makes the measurements is behind the brattice out of the air stream. The Pitot tube does not offer as disturbing an obstruction as the anemometer. A recording gauge may be employed to integrate fluctuations in air flow through that portion of the mine. No traverses are required because the reading may be at a single central point. An anemometer can be used with an orifice flow. The orifice will increase the air velocity at the measuring point, with correspondingly more accurate measurements where the normal air velocity through the passageway is low. Portable brattices might be devised with the cushioning rims which would seal against irregular rock surfaces where permanent rings were not available or feasible. The development by the Ventilation Committee of standard procedures and devices for the orifice measurement of the flow of mine ventilating air might be a desirable project for this coming year. C. M. Smith (author&apos;s reply)—Thank you for your discussion of my paper on underground anemometry. Your suggested method of measuring underground air flow is a novel one which might be applicable in some situations. It should be tested along with other suggested methods in any investigation of this subject. G. E. McElroy*—In spite of the adverse publicity that vane-anemometer methods of air measurement have had in the past and that contributed by the present paper, I endorse Mr. Erickovic&apos;s statement that anemometer traversing "has proved to be widely applicable, expeditious and simple" and add that available methods are accurate enough for the purposes for which they may be used. The fact that the great majority of minor mine officials assess relative changes in rates of air flow by comparison of crude vane-anemometer measurements, known to average 20 to 30 pct high, has no important bearing on this subject, because state inspection standards were based originally on such methods of air measurement. Federal inspection standards are based on actual rates of flow as determined by traversing, and interest in traversing methods is rapidly increasing. In considering traversing methods, three aspects are of major importance: (1) the absolute accuracy of calibrations; (2) the degree of interference with normal flow conditions introduced by traversing methods designed for accurate measurement by shaft-mounted instruments; and (3) the proper "method" factor to use for approximate measurements by hand-held instruments. With respect to absolute accuracy of calibrations, we have always placed reliance on calibrations made by the National Bureau of Standards, with which manufacturers&apos; calibrations have usually agreed very closely. It is therefore particularly disturbing to find7 that calibrations made previous to June 1947 are presumably about 5 pct in error because of excessive registry caused by the thin flat plates on which anemometers were mounted for calibration. Velocities corrected for calibration have therefore averaged about 5 pct low in all probability. In this connection, it is interesting to note that an anemometer calibrated against Pitot-tube measurement by a single-point method in the Bureau of Mines experimental coal nine in 1923 indicated this same difference of about 5 pct and that the same instrument calibrated by a traversing method in a metal mine some months later indicated a difference in the same direction of about 4 pct. These results are reported by the Bureau of Mines.8 Regarding the degree of interference, or changes in velocity distribution, caused by the position of the observer&apos;s body in traversing operations, misconceptions seem to be especially prevalent, resulting in increasing advocacy of methods, such as the "clear section" method outlined in this paper, that cause just the type of interferences that they are designed to avoid. The degree of interference for any method may be gauged easily by a few experiments with a velocity-pressure gauge connected to a Pitot tube or with an indicating velocity meter such as the Velometer. In an experiment cited by McElroy and Richardson,# a decrease of 5 pct was noted at ten widths upstream from a 6-in. plank, whereas an observer&apos;s body at about the maximum practical distance of 6 ft downstream from the instrument is only about four widths away. In the Bureau of Standards paper previously mentioned, it is recommended that supports used in calibrations be at least 16 widths downstream. In practice, therefore, a downstream position of the observer is ruled out as far as accurate measurement is concerned. Operation of the anemometer by rigid shaft support from a point outside the section is seldom practicable; however, accurate results can be obtained, with the anemometer rigidly attached to a short shaft and held at arm&apos;s length, by an observer advancing across the traversed section while he faces the opposite wall and stands sideways to the current, provided that he keeps the instrument at least 3 ft away from his body at all times and traverses the entire section with it. If the traverse can be started with the observer in a side recess, the entire section can be covered in one operation. Normally, it would be covered in two half-sections. The presence of the observer&apos;s body does not, as is commonly supposed, increase the average velocity throughout the remaining part of the section. Rather, the velocities 1 to 2 ft on either side of his body are increased, but the distribution of velocities throughout the rest of the cross section remains normal, and a traverse made as stated gives a true average velocity for normal-flow conditions. Regarding the proper "method" factor for accounting for interference in the approximate methods of traversing with hand-held instruments, here again confusion prevails, for which the writer must assume some of the blame. Comparison of consecutive traverses made by shaft-held and hand-held 4-in. anemometers in field work after the tests reported by McElroy and Richardson&apos; gave method factors

Jan 1, 1950

By H. L. Gilbert, W. W. Stephens

Production of anhydrous zirconium tetrachloride by direct chlor- ination of a zirconium oxide carbon mixture in a silica-brick-lined chlorinator is described. Theory and thermodynamics of reactions are discussed. A pilot-model chlorinator and full-scale production equipment are described and operating data are included. ANHYDROUS zirconium tetrachloride required as a starting material in the Kroll process for production of ductile zirconium has been produced in this country principally by chlorination of the carbide or "carbonitride" made by reduction of zircon sand concentrates with carbon in the arc furnace. This process and the equipment used have been fully described.&apos;.&apos; It is well known that zircon or badde-leyite ores may be chlorinated directly with chlorine in the presence of carbon, and this one-step approach would seem at first glance to be preferable to the two steps involved in the carbide-chloride operation. As previously discussed1 considerations leading to adoption of the longer process were: 1—Chlorination of the carbide proceeds rapidly at temperatures below 500°C, whereas temperatures above 900°C were considered necessary for chlorination of zircon-carbon mixtures. 2—The highly exothermic nature of the carbide-chlorine reaction makes it self-sustaining, while heat must be supplied continuously in direct chlorination of the ore. 3—Silicon was thought to be chlorinated along with zirconium in the direct chlorination of zircon, leading to high chlorine consumption. In production of carbide in the arc furnace, silicon is driven off as silicon monoxide and does not enter the chlorinator. 4—For efficient direct chlorination, the ore must be finely ground and intimately mixed with carbon, and the mixture briquetted. 5—Direct chlorination requires a much larger chlorinator for a given production capacity than chlorination of carbide. Recently it has become necessary to produce large quantities of zirconium metal from a chemically purified zirconium oxide. Since the cost of the latter is high, the relatively high losses encountered in production of carbide in the arc furnace could not be tolerated, and a graphite resistor furnace5 was developed for production of carbide, which was then chlorinated in equipment previously used for chlorinating arc-furnace carbide. This method of operation was quite satisfactory, and losses in the carbid-ing step were minimized. However, operating costs were relatively high, and the process did not lend itself particularly well to large-scale operation because of the multiplicity of small units required and the hand labor needed to load, unload, and maintain the furnaces. To chlorinate the carbide, a vertical-shaft chlorinator was used in which the charge was heated with a central split graphite-rod resistor.&apos; Operation of this chlorinator was somewhat less satisfactory with the resistor furnace carbide than it had been with the arc-furnace carbide due, principally, to the differences in physical properties of the carbide. The arc-furnace carbide is obtained as a fused metallic-appearing mass which can be crushed to -1/4 in. with production of a minimum of fines, whereas the resistor-furnace product is lightly sintered and produces a large proportion of fines in crushing and handling. These fines tend to pack in the chlorinator and promote channeling. which results in poor chlorine efficiency and low capacity. Data based on production of 22,000 lb of chloride from resistor-furnace carbide in this equipment are shown in Table I. Necessity for increasing production of chloride from about 2,000 to 15,000 lb per week led to further investigation of possible methods for direct chlorination of the oxide or oxide-carbon mixtures. Direct chlorination of the pure oxide presents much less difficulty than chlorination of zircon sand or oxide ores. The oxide is easily ground to —200 mesh, and no silica is present to cause excessive consumption of chlorine or contamination of the product. Theoretical Considerations The principal chemical reactions involved in chlorination of mixtures of zirconium oxide and carbon may be written as follows: ½ ZrO2 (c) + C(c) + Cl2(g) = ½ ZrCl,(g) + CO(g) [I] 1/2 ZrO, (c) + CO(g) + Cl2(g) = ½ ZrCl,(g) + CO2(g) 121 ½ ZrO2 (c) + ½ C(C) + Cl2(g) = 1/2 ZrCl4(g) + ½ CO2(g) [3]

Jan 1, 1953

By J. Szekely, Robert G. Lee

Calculations are presented for the prediction of the combined radiative-convective heat loss from molten steel held in a ladle, covered by initially molten slag. A mathematical formulation is given and numerical solutions are presented of the resultant partial differential equation. It is shown that a slag layer of about 2 to 3 in. thick would prevent any significant heat loss from the top metal surface for Periods of up to 1 hr. For thinner slag layers (or in the absence of the protective slag) appreciable heat losses would occur; plots given in the paper allow the quantitative evaluation of the fall in the metal temperature due to this heat loss. The ironmaking-steel processing operation involves several intermediate steps where the molten metal is held or transfered in ladles. The quantitative assessment of the heat losses occurring under these conditions is of considerable importance in process design considerations since the temperature of the metal may have to meet rigid specifications in any given part of the processing sequence. While the metal is held in a ladle, heat is lost by two mechanisms: i) by conduction into the ladle walls; ii) by radiation and convection from the top surface. Calculations relating to the conductive heat loss, are readily made and information may be found in the literature both on data pertaining to metallurgical situations and on the techniques that are available for performing additional computations.&apos;-3 The evaluation of the combined convective-radiative heat loss is less straightforward, especially when there exists a protective slag layer covering the metal. In this latter case, as the top slag surface loses heat partial solidification may occur and the latent heat thus given up by the slag may represent an effective barrier to the heat loss from the metal. A semiquantitative assessment of the role played by the slag in preventing heat loss from the metal has been made in a previous paper,4 where it was shown that a slag layer 6 in. thick would act as a near-perfect insulator for periods of up to 2 hr. However, this first paper reported on an essentially preliminary investigation that considered only one slag layer thickness; furthermore, the boundary conditions used for the calculations were somewhat restrictive since no allowance was made for changes in the metal temperature. The purpose of this second paper is to extend the scope of the preliminary investigation previously reported by: i) considering several slag layer thicknesses; ii) allowing for variations in the metal temperature; and finally iii) by performing calculations on the net loss from the metal rather than from the slag surface. FORMULATION Consider a slag phase extending from y = 0 to y = L1, covering a metal phase that occupies the region extending from y = Ll to y = L, as illustrated in Fig. 1. Denote the slag and metal temperatures by TS and Tm and assume that initially both slag and metal are molten, having a uniform temperature Ti. At time = zero the surface represented by the y = 0 plane (top slag surface) is brought into contact with cold air, the temperature of which is given as T Thus for time > zero, heat will be transfered from the slag to the air by natural convection and thermal radiation; as a result of this heat loss the slag temperature will fall and after some time a solid phase is formed and a solidification boundary (more realistically described as a zone) will move progressively from y = 0 toward y = L1. Once the region of temperature gradients, moving ahead of the solidification zone, reach the slag-metal interface (y = L1) there will be a net transfer of heat from the metal through the slag to the cold air, across the y = 0 plane. In the formulation of moving boundary problems, involving phase changes that require computer solution, it is convenient to represent both phases by a single equation, making allowance for the latent heat of solidification by assigning an appropriate temperature dependence to the specific heat content. Thus the energy equation for the whole of the slag may be written as follows:

Jan 1, 1969

By Ian MacGregor

INTRODUCTION I retired in 1977 and have been enjoying myself ever since doing odd jobs for various people - on occasions, politicians. I commend it to you as post retirement - get involved in situations that nobody else will take on. It produces all sorts of excitement, it is a lot of fun and, courtesy of the media, you provide great raw material for their current Punch and Judy shows. The example of the adventures in the Escondida financing made me think about the history of the minerals industry in the United States. In order to see more clearly what is going to happen in the future, we have to look at the past and what we have learned, and try to extrapolate from that a scenario for the future. This paper is about global material cycles and the financial needs of the extractive industries. Since about 2,000 years B.C. we have been living in the Iron Age. Since about 4,000 B.C. we have been living in the Copper Age. These have been the two foundations of modern industrial society. Understanding the use of metals and materials helps us understand the civilization in which we live today. Let me focus more closely on the United States, particularly in this century. In the first two-thirds of this century the United States was growing and expanding and developing. It was providing for its inhabitants a society which was enjoying the benefits of a rising standard of living, improved opportunities and the ability to support a population growing at quite a dramatic rate. It is still growing. In the Department of Commerce, there is a clock that shows how many people there are in the United States. It turns all the time; the numbers are going up. On January 21,1980, when Ronald Reagan was sworn in as President, the number on that clock was a little over 220 million, and in the last week before he left, it was travelling towards 245 million. During his period in Washington, we have added , the equivalent of one total state such as California. That gives the dimensions of the U.S. appetite for raw material inputs. The inputs are of critical importance. COMPARATIVE ADVANTAGE During the first two-thirds of this century, the minerals sector enjoyed comparative advantages. There was a clear understanding of the importance of minerals enshrined in the tax legislation of the country from the early decades of this century - the depletion allowance. No other country identified as clearly that the supplies of minerals were finite. As they are used up, just as with a wearing-out machine, one must get capital back to replace the resource when the time comes. Within the last decade, that idea has been questioned and criticised. It is now seen as giving money away to a preferred part of the economy, instead of recognizing what depletion is all about. In previous times it insured that the U.S. always was able to regenerate the capital in its minerals industry, so that it could remain competitive. The U.S. had plenty of everything, including energy. We had a growing population. Equally important, we had an agriculture that led the world in productivity, even as far back as the turn of the century. It continues to do so. We still retain that comparative advantage, an ability to produce our foodstuffs in quantity and quality at excellent cost. Look at another ingredient of our comparative advantage, the energy position. We ran through the first two-thirds of the century with ever- decreasing energy costs. In the 1940s and 1950s the coal industry went into a decline. Why? Because the $3.00 - $5.00 a ton cost of mining coal in the U.S. represented relatively expensive energy. A ton of good grade bituminous coal contains 25 million BTU&apos;s. $0.20 a million BTU&apos;s represents $5.00 a ton at the pithead. That price was undermined by the increasing ability to find petroleum at even lower costs. At the apex of the growth of development of the American dominated

Jan 1, 1990

By J. R. Jargon, H. K. van Poollen

Non-Newtonian fluids may be injected into a reservoir during secondary recovery operations. The non-Newtonian fluid used in this work is a power-law type of fluid that is, the viscosity of the fluid decreases as the flow rate or rate of shear increases. This paper presents equations for steady-state linear and radial flow of such fluids, transient behavior results from a finite difference model of a radial system, and transient behavior results from a field test. The equations that describe the flow of a non-Newtonian fluid are non-linear and are solved numerically. Finite difference solutions are presented as curves of dimensionless pressure drop at the wellbore us dimensionless time for a constant injection rate. Solutions were obtained for 5-percent, 10-percent and 100-percent PV of a non-Newtonian fluid for injection rates of 1, 10, 100 and 1000 cc/sec and for a 5 percent PV of non-Newtonian fluid located at T = Tw&apos; 3, 10, 20, 50 and 100 ft for a flow rate of 1 cc/sec. The buildup curves do not exhibit a stmight-line portion as is the case for Newtonian flow through porous media. Correlations also are shown for the productivity index vs rate for the computer model study and the field tests. INTRODUCTION During various secondary recovery operations non-Newtonian fluids are injected. Such fluids, in general, include thickened water and gelled fluids. The term "non-Newtonian" implies that the viscosity is not only dependent upon temperature and pressure, but also on the rate of shear that is applied to the fluid. For example, water, which is a Newtonian fluid, will have essentially the same viscosity no matter what rate of shear is applied. A pseudo-plastic fluid exhibits a decreasing viscosity when higher rates of shear are applied; a dilated fluid has an increasing viscosity with increasing rates of shear. The objective of studies performed and described in this report is to obtain relationships and mathematical and empirical descriptions of the flow of non-Newtonian fluids through porous media. Simple mathematical relationships, computer studies that include the unsteady-state behavior of such fluids, and field tests are used. PREVIOUS LABORATORY STUDIES Laboratory studies have been performed by several investigators. The fluids used in one investigationl were surfactant-stabilized dispersions of water in hydrocarbons. Its Fig. 6 is reproduced as our Fig. 1. This effective-viscosity vs frontal-velocity diagram shows that this fluid is of the power-law type over a fairly large range of frontal velocities. Gogartyl developed an equation for the effective viscosity as a function of shear rate that reduces to Eq. 2 for frontal velocities greater than about 2 ft/D. The fluid characteristics used in the present study are similar to those reported by Gogarty. Christopher and Middleman2 performed experiments with power-law fluids flowing through pipes packed with glass beads. They used a modified Blake-Koteny

Jan 1, 1970

By F. V. Lenel

Two years ago in Chicago a seminar was held on the theory of sintering of pure metal powders. As an introduction to this seminar D,.. Rhines1 gave an excellent survey of the literature on this subject, His method of presentatioll was to summarize the experimental observations on sintering and from them to develop a composite theory of sintering in which he combined all the important contributions to the theory into one organic whole. In contrast to the mechanism of sintering of pure metal powders, which, of course, al\vays takes place in the solid phase, sintering j,, the presence of a liquid phase, cannot be treated as one unified mechanism, such as the sintering mechanism of pure metals. The reason is that there are really several mechanisms depending upon the type of alloy system which is sintered and the field of its con. stitutional diagram ill which the sintering takes place, Two mechanisms have heen investigated and will be reviewed in this paper. Ill the first lnechanism the liquid is presellt during the entire time while the compacts are at the sintering temperature, In other words they are sintered between the liquidus and the solidus of the alloy system and are ~e~erogelleous during the entire sintering cycle. The second mechallism applies to alloys in which the liquid phase is formed during the sintering process, but disappears before the sintering process is completed through diffusion and formation of a solid solution. These alloys are therefore homogeneous at the end of the sintering cycle. This second mechanism is more complicated because of the two stages of sintering, the liquid and the solid, as they may be called. It is therefore not surprising that the few detailed investigations of sintering in the presence of a liquid phase have been concerned with heterogeneous sintered alloys, namely the tungsten-nickel-sinter ing which is the so-called "heavy alloy," and the cemented carbides. The mechanism of sintering of these systems is characterized by the fact that theoretical or near theoretical density is attained during sintering, while simultaneously a distinct grain growth takes place through solution of the smallest grains of the solid phase in the liquid phase and reprecipitation on the larger grains. In the first part of the survey this sintering mechanism which is called for short the heavy alloy mechanism is treated in detail. In the second part Of the survey experimental observations for the sintering of homogeneous sintered alloys where the liquid phase disappears before the completion of sintering are discussed. Much less systematic work has been done on these alloys to which the conlmercially important Porous bronzes and the iron-nickel-aluminum permanent magnet alloys belong. Emphasis will be laid throughout this survey upon the microstructural and density changes during sintcring because the Sintering mechanism can usually be described directly by these changes. Other

Jan 1, 1949

By Fred C. Bond

COMPRESSION TESTS THE standard method of determining the crushing resistance of rocks consists of crushing prepared shapes under slow compression, and expressing the ultimate crushing resistance at the load causing failure in pounds per square inch of crossectional area perpendicular to the crushing force, with the height approximately equal to the diameter. One-inch cubes or cylindrical drill cores are commonly used. Cubes and drill cores are cut by a high-speed steel disk with diamond dust embedded in the edges, and running under a water spray. Such a saw will cut one square inch or more of hard stone per minute. An oiled, spherical, swivel compression block of small diameter should be used to equalize the pressure. Each sample is crushed between sheets of cardboard or blotting paper extending about one-half inch beyond the edges of the stone. A cloth is wrapped around it to prevent dangerous explosive shattering. It is probably preferable to place the swivel block below the specimen, with either a rigid head or another swivel block above. Whatever the arrangement, the conditions should be selected that will develop the maximum crushing strength of the specimen, even though its top and bottom planes may not be precisely parallel. The cardboard pads may be dispensed with if the ends of the specimens are polished smooth. The use of lead pads was discontinued because of the observed tendency to split the material. The samples are weighed and measured before breaking, and the density is calculated. Cubes that contain bedding planes or veinlets are usually placed with these vertical and parallel to the crushing force, in order to develop the maximum resistance possible. A number of compression tests were made on Canadian ores, by Forrest Nagler, of the Allis-Chalmers Manufacturing Co., and are listed alphabetically in Table 1. These tests were made on one inch cubes or on diamond-drill cores approximately one inch in diameter and one inch long. The drill cores represent virgin rock, while the cubes may have been previously weakened by explosives. Compression tests have been made in the Allis-Chalmers Milwaukee Laboratory on a number of ores and other materials. These are listed alphabetically in Table 2. In test No. 1339, on limestone from the Hanna Coal Co., comparison was made between 2 by 2-in. diamond-drill core segments and one-inch cubes cut from the cores in the same strata. Results for the cylinders are tabulated directly below those for the cubes. The bedding planes were normal to the compressive force in breaking the cylinders and parallel in breaking the cubes, which accounts for the lower strength obtained from the cylinders. The percentage of net linear compression at the ultimate pressure was measured in test No. 1055 for Fontana TVA quartzite as 0.30 for cube A, and 0.37 for cube D.

Jan 1, 1946

By D. W. Mitchell

While pure magnesium does not corrode rapidly the presence of even very small quantities of certain other metals accelerates corrosion remarkably. Because magnesium is such an electropositive metal (E° = +2.34 volts)&apos; the presence of substances that permit galvanic activity is extremely undesirable. It is stated that as little as 0.017 pct of iron in magnesium causes a marked increase in the corrosion rate.2 Magnesium and its alloys are commonly melted in iron or steel pots; consequently the commercial metal has the opportunity to pick up iron if it has that tendency. Quantitative data on the solubility of iron in magnesium are of practical as well as of academic interest. Some magnesium alloy castings tend to have large grain size unless suitable treatment is given to the liquid metal just prior to casting. One method of grain refinement consists of "superheating" the metal just prior to casting. "Superhcating" is heating to a temperature 100 to 200°C above the casting temperature and holding for a short period of time. Superheated metal has a much finer grain than unsuperheated metal. A satisfactory explanation of this phenomenon has not been forthcoming. One hypothesis attributes the grain refinement in the superheated metal to precipitation of some phase that is mis-cible at the temperature of superheat but not miscible just below the casting temperature and which is thought to provide nucleating centers for crystallization. This phase, if the foregoing hypothesis is correct, would probably be a metal or an intermetal-lic compound. Of the metals likely to be found in magnesium alloys, iron is one of the most likely candidates for the role of grain refiner for iron occurs in small quantities in all commercial magnesium and magnesium alloys. A. Beck³ states that magnesium, in either the liquid or the solid state, cannot dissolve iron and that iron exists as a finely divided suspension which does not settle out because of its extreme fineness. Beck also presents a micrograph showing a large iron dendrite in magnesium but does not comment on the crystallization of iron from a solution supposed to contain none. It is doubtful that dendrite formation from suspended iron -would be possible. This suspension theory is the basis of patents4 covering the removal of iron from magnesium by allowing the iron particles to act as nuclei for the formation of primary magnesium crystals which can be separated by segregation. Also covered by patents is the removal of iron by the addition of another metal to form an intermetallic compound with the iron which, because of its increased particle size, is able to settle out of the molten metal. Here, too, quantitative information on the extent of iron solubility in liquid magnesium is of more than theoretical interest. M. Hansen,5 in his critical collection of the data on binary alloys, states that iron is not soluble in liquid magnesium. W. R. D. Jones% ho studied the physical properties of low iron magnesium alloys observed a 57°

Jan 1, 1949

By David W. Mitchell

WHILE pure magnesium does not corrode rapidly the presence of even very small quantities of certain other metals accelerates corrosion remarkably. Because magnesium is such an electropositive metal (E° = +2.34 volts)&apos; the presence of substances that permit galvanic activity is extremely undesirable. It is stated that as little as 0.017 pct of iron in magnesium causes a marked increase in the corrosion rate.2 Magnesium and its alloys are commonly melted in iron or steel pots; consequently the commercial metal has the opportunity to pick up iron if it has that tendency. Quantitative data on the solubility of iron in magnesium are of practical as well as of academic interest. Some magnesium alloy castings tend to have large grain size unless suitable treatment is given to the liquid metal just prior to casting. One method of grain refinement consists of "superheating" the metal just prior to casting. "Superheating" is heating to a temperature 100 to 200°C above the casting temperature and holding for a short period of time. Superheated metal has a much finer grain than unsuperheated metal. A satisfactory explanation of this phenomenon has not been forthcoming. One hypothesis attributes the grain refinement in the superheated metal to precipitation of some phase that is miscible at the temperature of superheat but not miscible just below the casting temperature and which is thought to provide nucleating centers for crystallization. This phase, if the foregoing hypothesis is correct, would probably be a metal or an intermetallic compound. Of the metals likely to be found in magnesium alloys, iron is one of the most likely candidates for the role of grain refiner for iron occurs in small quantities in all commercial magnesium and magnesium alloys. A. Beck&apos; states that magnesium, in either the liquid or the solid state, cannot dissolve iron and that iron exists as a finely divided suspension which does not settle out because of its extreme fineness. Beck also presents a micrograph showing a large iron dendrite in magnesium but does not comment on the crystallization of iron from a solution supposed to contain none. It is doubtful that dendrite formation from suspended iron would be possible. This suspension theory is the basis of patents4 covering the removal of iron from magnesium by allowing the iron particles to act as nuclei for the formation of primary magnesium crystals which can be separated by segregation. Also covered by patents is the removal of iron by the addition of another metal to form an intermetallic compound with the iron which, because of its increased particle size, is able to settle out of the molten metal. Here, too, quantitative information on the extent of iron solubility in liquid magnesium is of more than theoretical interest. M. Hansen,5 in his critical collection of the data on binary alloys, states that iron is not soluble in liquid magnesium. W. R. D. Jones6 who studied the physical properties of low iron magnesium alloys observed a

Jan 1, 1948

By T. Glen Heslop, Dennis H. Laubscher

INTRODUCTION In all cave mining operations waste is drawn mixed in the ore, hauled, hoisted and milled. This waste usually contains some mineral which partly covers some of the costs of mining and milling it, but it still re- presents a loss of profits as the production rate from the mine is matched to the mill capacity - every tonne of waste milled replaces a tonne of profitable ore. Draw control aims to reduce the amount of waste drawn and to increase the amount of ore recovered. It is also used to regulate the grade of ore flowing to the mill and in block cave blocks to prevent potential stress build- up on the major apexes above grizzly drifts. This paper summarises what has been learnt about draw control in the forty years that cave mining methods have been used in the three largest underground chrysotile asbestos mines in the southern hemisphere. The four mining methods to which draw control has been applied are: a. Block caving. b. Shrinkage with overburden caving. c. Sublevel caving. d. Open stoping. Open stoping has been included because on a number of occasions severe falls from the back and side of these open stopes have occurred. Without full records of where and when ore is blasted, when and how much ore has been drawn from each of the drawpoints, it would be impossible to analyse the migration of ore and fallen waste in the stope. Several problems peculiar to chrysotile asbestos ore bodies affect the practice of draw control. Heterogeneous fragmentation, cavities, side pressures and the interaction of stresses in the caved ore affect the migration of ore under draw. Fibre is normally not developed uniformly throughout the ore bodies, values vary from place to place but there is usually a general decrease upwards and the hangingwall contact is an economic limit. The immediate hangingwall rock frequently looks like the ore body hostrock and is difficult to recognise in a drawpoint, and cannot be used as the marker for stopping draw. During caving or shrinkage, a large proportion of fibre becomes detached from the host rock and re- ports in the fines, which migrate differently to the larger rocks. The fines are not sticky, and flow easily if loose and dry, but compacted, the fibre and rock fines con- solidate the ore. During draw the proportion of fines in the ore varies considerably, as fines either flow relatively freely between the large rocks or become trapped and compacted above large rocks or hangups. This creates a large sampling error in drawpoint grade assessments, whether done visually or by bulk sampling, unless the assessments are done frequently. The method of draw control used is to regulate the draw rate in an attempt to lower the ore-waste interface as evenly as possible, and to finally stop working the drawpoint when a calculated tonnage has been drawn.

Jan 1, 1981

Using 5-p Ni powder and 0.018 ,u A1203, oxide dispersion strengthened nickel alloys were prepared by mechanical mixing of powders, followed by compaction, sintering and extrusion. Processing variables such as cooling of the powder during mixing, electrostatic discharging, mixing time, amount of oxide, and so forth, were studied to determine the effects on room-temperature tensile and high-temperature creep-rupture properties and on the reproducibility of alloys. MECHANICAL mixing of metal and metal oxide powders provides one of the simplest and most inexpensive methods for preparing dispersion strengthened metal-metal oxide alloys;" however, because of large differences in specific gravity, powder particle diameters and other variables, difficulties have been encountered in reproducing alloys and alloy properties. Using 5-p Ni powder and 0.018 p A1203, the following variables were investigated: a) Blending techniques b) Blending time c) Volume fraction of oxide for the selected pow- ders d) Sintering of compacts before extrusion Cooling of Powders During Blending— Cremens and Grant had indicated that dry ball milling of the powder mixture gave superior results over other procedures. When ultra-fine metal and oxide powders became available and comminution was unnecessary or offered little gain, the use of dry mixing in a Waring Blendor gave optimum results.&apos; In the current study, observations of the powder mixtures indicated that considerable heat was generated due to the high speed of the Waring Blendor (15,000 rpm). This heating appeared to produce some degree of clustering and pelletizing. To minimize such agglomeration, a Pyrex cooling coil was installed in the mixing vessel. A double coil, water cooled, proved to be quite satisfactory, providing a fair degree of cooling along with good erosion resistance. Discharge of Electrostatic Charge—Electron-microscopy showed that agglomeration was minimized by the cooling procedure but was not entirely eliminated. It appeared that some of the ag- glomeration had to be explained on the basis of an electrostatic attraction due to the high-speed blending operation. These charges had the effect of promoting the growth of alumina clusters. At the end of the blending procedure such clusters had an effective size which was as much as 50 times that of the individual powder particles and were very tightly bunched. Several methods were utilized to try to discharge the electrostatic forces to minimize the clustering. One successful method, utilized in alloy 17, involved continuous exposure during blending to a radium source of 50 millicuries at a distance of 1.5 ft from the container. A number of other methods were noted to work equally well. Fig. 1 shows the more open oxide structure which results from such a treatment, permitting the matrix metal to penetrate into the cluster more effectively. The combination of powder cooling and discharging led to noticeably improved structures. Fig. 2 shows a cross and a longitudinal section in alloy M-17, one of the best alloys, illustrating the marked decrease

Jan 1, 1961